Functionalized poly(vinylidene fluoride) nanohybrid for superior fuel cell membrane

Author's Accepted Manuscript

Functionalized poly(vinylidene fluoride) nanohybrid for superior fuel cell membrane Karun Kumar Jana, Chumki Charan, Vinod K. Shahi, Kheyanath Mitra, Biswajit Ray, Dipak Rana, Pralay Maiti

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PII: DOI: Reference:

S0376-7388(15)00092-7 http://dx.doi.org/10.1016/j.memsci.2015.01.053 MEMSCI13457

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Journal of Membrane Science

Received date: 8 November 2014 Revised date: 20 December 2014 Accepted date: 22 January 2015 Cite this article as: Karun Kumar Jana, Chumki Charan, Vinod K. Shahi, Kheyanath Mitra, Biswajit Ray, Dipak Rana, Pralay Maiti, Functionalized poly (vinylidene fluoride) nanohybrid for superior fuel cell membrane, Journal of Membrane Science, http://dx.doi.org/10.1016/j.memsci.2015.01.053 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Functionalized Poly(vinylidene fluoride) Nanohybrid for Superior Fuel Cell Membrane

Karun Kumar Jana1, Chumki Charan2, Vinod K. Shahi2, Kheyanath Mitra3, Biswajit Ray3, Dipak Rana4, and Pralay Maiti

1

*1

School of Materials Science and Technology, Indian Institute of Technology (Banaras

Hindu University), Varanasi 221 005, India 2

Electro-Membrane Processes Division, Central Salt and Marine Chemicals Research

Institute, Bhavnagar 364002, Gujarat, India 3

Department of Chemistry, Banaras Hindu University, Varanasi 221 005, India

4

Industrial Membrane Research Institute, Department of Chemical and Biological

Engineering, University of Ottawa, 161 Louis Pasteur St., Ottawa, ON, KIN 6N5, Canada

*

Correspondence should be made to [email protected] (P. Maiti)

1

Abstract Functionalization of poly(vinylidene fluoride) (PVDF) nanohybrid has been performed in template system using two-dimensional layered silicate and superior fuel cell membrane has been demonstrated. Sulfonation of nanohybrid has been carried out at control condition to maintain the mechanical stiffness and toughness of the membrane using chlorosulfonic acid and the results have been compared with pure PVDF. The sulfonation and its relative extent have been confirmed through NMR, FTIR and UV-Vis measurements showing greater degree of functionalization in nanohybrid which arises from the specific arrangement of polymer chains on top of nanoplatelets. The structural change over from common crystallized form α− to piezoelectric β−phase in nanohybrid has been established and the amount of β−phase has been enhanced after sulfonation as evident from deconvoluted XRD patterns and DSC measurement. A plausible mechanism has been proposed for this improvement which led to the formation of smart membrane. Essential criteria of an ideal membrane have been verified through high water uptake, low permeability and hydrophilic nature by measuring contact angle. The molecular level clustering due to the attachment of sulfonate group in main chain has been explored which in turn explain the higher barrier property both for gas and liquid (fuel). Proton conductivity of functionalized nanohybrid has been found to be quite high along with significantly low methanol cross over as compared to standard Nafion membrane. I-V characteristics of the nanohybrid membrane show high potential at low current density with considerably lower value of slope. Membrane electrode assembly using functionalized nanohybrid exhibit significantly high value of current density and prove its worth for superior fuel cell membrane using common thermoplastic polymer. Keywords: PVDF - nanohybrid, sulfonation, proton exchange membrane, Structure 2

Introduction Renewable energy sources and the related technologies have enormous challenge in recent time as different fossil fuels are being used up and great demand of energy in portable power applications [1]. Amongst them proton exchange membrane or polymer electrolyte membrane (PEM) is one of the most promising components due to its high power density in energy conversion devices [2]. PEM fuel cell uses a solid polymer membrane (a thin plastic film) and act as both electrolyte and separator [3] The function of membrane, generally prepared from respective ionomers (sulfonate, phosphonate and chlorinate ions, etc.), is to transport proton from anode to cathode but should not conduct electrons [4]. Nafion (ionomer with a fluoropolymer) is the most common membrane with good proton conductivity (σ ≥ 10-2 s cm-1) in low temperature was developed by DuPont Inc [5,6]. However, it has several drawbacks such as high production cost, high methanol crossover and significant decrease of its conductivity at high temperature especially above 80 oC [7]. In recent decades, many researchers have developed new functionalized membranes to overcome these problems using various techniques, such as graft copolymerization [8], block copolymerization [9], radiation [10] and sol gel methods [11]. Recently, Ameduri et al. have shown a proton conducting electrolyte membrane based on fluoropolymer which has achieved proton conductivity up to 9 m.S cm-1[12]. The basic requirements of a proton exchange membrane are low fuel permeability, good proton transport properties, high mechanical and chemical stability, high proton conductivity together with low production cost [13-15]. It is well established that the super hydrophobic, mechanically tough and electrochemically stable insulating poly(vinylidene fluoride) (PVDF) is the most promising polymer as far cost is concerned as compared to Nafion for fuel cell membrane applications [16,17]. 3

PVDF can crystallizes in five molecular configurations such as α, β, γ, δ and ε phases ¹

depending on the geometric conformations like nonpolar ( TGT G ), polar (TTTT), partially ¹

¹

¹

polar (T3GT3 G ), partially polar ( TGT G ) and nonpolar ( (T3GT3 G ) phases, respectively [18]. Amongst them, the polar β-phase (all trans) is responsible for piezo- and pyroelectric properties and several applications have been originated using this form of PVDF such as sensors and electromechanical systems [19]. Usually, 25-150 µm thickness membrane is the good compromise between ionic strength, gas permeability and mechanical strength [20]. Incorporation of some organic or inorganic filler into the polymer matrix, so called nanocomposites, has shown tremendous improvement in various properties including mechanical, thermal, physical and chemical due to their reinforcing effect [21-24]. In this work, our interest is focused on the preparation of a membrane with high proton conductivity with lower cost through direct sulfonation of poly(vinylidene fluoride) (PVDF) in presence of nanoparticle (template system). Nonconducting poly(vinylidene fluoride) (PVDF) has been functionalized in presence of organically modified nanoclay to prepare the ionomer for PEM fuel cell applications. The effect of nanoparticles has been established for increasing the rate of functionalization which in turn improves the mechanical along with electrical properties, the basis for using membrane for practical applications. Other properties required for any membrane, like water uptake, hydrophilicity, gas permeation are measured to check the suitability of this novel material as fuel cell membrane. In addition, the molecular level phenomena associated with functionalization and details of structural characterization have been explored to find the cause behind the paradigm improvement in properties using a common material which has proven to be better material for fuel cell membrane. Experimental Section Materials: The polymer used in this study was poly(vinylidene fluoride) (PVDF) (SOLEF 6008; Ausimont, Italy), with a melt flow index of 24 g/10 min at 230 oC under a 5kg load. Organically modified nanoclay was Cloisite 30B [bis(hydroxyethyl) methyl tallow ammonium ion exchanged montmorillonite], Southern clay products Inc. (Gonzales, TX), 4

which was used as the nanofiller (fine powder). Tallow is a mixture of C16 and C18 long chain alkenes. Sulfonation of the samples was performed using chlorosulphonic acid (LOBA Chemie) for application point of view. Nanohybrids and film preparation: A home-built chip sizer was used for converting PVDF beads into powder. Required amounts of PVDF powder and nanoclay (4 wt.%) were first mixed thoroughly in a high speed (1000 rpm) mixer before putting into the extruder. Melt extrusion technique has been used for nanohybrid preparation using a twinscrew extruder (Hakke Mini Lab) at 205 oC for 10 min at a high shear rate of 100 rpm. Henceforth, the PVDF nanohybrid with 4 wt.% of nanoclay in the polymer matrix will be used as “NH”. After the extrusion process, both the pure PVDF and nanohybrid (NH) were melt-pressed into a thin film of ~130 µm thick of size 4×4 cm2 in a compression molding machine (S. D. Scientific Ltd.) at 190 oC under 5 tons of pressure for functionalization. Preparation of the membrane through functionalization: The 130 µm thick films of PVDF and its nanohybrids (NH) were functionalized directly using chlorosulphonic acid at different time and temperature. The sulfonation reactions were allowed to proceed for varying time (3-15 hrs) at 65 oC and over a wide temperature range of 40 to 100 oC for a fixed 4 h time to understand the effect of functionalization time and temperature. Finally, an optimized condition of temperature and time was chosen so that mechanically stable polymer film was produced after chemical modification for their ready use. The sulfonated films were washed with deionized water until the residual water had a pH of 7 and remaining water at the surface was absorbed with a filter paper. The films were then dried at 60 oC for 24 h under reduced pressure. Henceforth, we will term the sulfonated species as PVDF-s or NH-s for PVDF and nanohybrid, respectively. Separate set of samples were 5

used for various characterization. To evaluate the water uptake (Wwater) at 25 oC, the membranes (4×4 cm2) were first dried under vacuum at 60 oC for 24 h to obtain their dry weight (Wdry). Then, the membranes were placed in deionized water for 5 h to calculate the weights of the water-swollen (Wwet) specimens. The water uptake was measured following the eqn. 1 and has been taken as the average of three samples.

Wwater =

Wwet − Wdry

(1) × 100% Wdry To measure the ion exchange capacity (IEC) of the membranes, firstly the membrane samples were immersed into 2 M HCl to convert the sulfonated species into acidic form. The membranes were washed with double distilled water to remove the adsorbed acid. 2 M NaCl solution was used for equilibrating the membrane for 24 h for ion exchange to take place. The remaining solution was titrated with 0.025 M NaOH solution using phenolphthalein as an indicator. The IEC values were calculated following the eqn. 2;

VNaOH S NaOH (2) Wdry are the volume and molar concentration of NaOH solution used IEC =

where, VNaOH and SNaOH

for titration. Wdry is the weight of dried samples.

Molecular analyses: The degree of sulfonation of the membranes were assessed through 1

H–NMR spectra using a JEOL AL300 spectrometer at room temperature in deuterated

dimethyl sulfoxide (DMSO-d6) as solvent and are reported in parts per million (δ). The degree of sulfonation DS (%) was measured from the ratio of integrals areas of the peaks, following the eqn. 3;

6

M ×100 (3) M+N where, M and N are the integration areas of the peaks assigned as c, d, e and a, a/, b as DS ( % ) =

mentioned in Fig. 1a. The crystalline phase of the films were determined using a Thermo Scientific FTIR (NICOLET-6700) in the ATR mode in the wavenumber range of 650-4000 cm-1. Each spectrum was recorded by accumulating 100 scans with a peak resolution of 4 cm-1. To understand the functionalization, UV-Vis spectra (Shimadzu UV-17001) were recorded in the reflection mode using the solvent cast films on quartz plates in the wave length range of 200-1100 nm. Gel permeation chromatography (GPC; Younglin ACME 9000) of pre- and post- functionalized samples was conducted at 70 oC in DMF as eluent at a flow rate of 0.5 ml/min against polystyrene standards. Gas permeation: Gas permeability of the membranes (pre- and post- functionalized) was measured at a constant pressure using three parallel cells method [25]. A rectangle film of 5 cm2 area and 0.9 mm thick membrane was placed in the cell and was sealed using aluminum foil, glue and cello tape. Feed pressure was set at 80 psi while the permeate side was maintained at the atmospheric pressure. Soap bubble flow meter was used for the determination of gas permeation rate. Every experiment was carried out three times for reproducibility and the average of those results is reported. The gas permeability was determined from the following eqn. 4: 10 ml (4) tmin × 60 is the time (sec) of gas flow rate of the permeate gas passing through the

Gas flow rate (Gfr) =

where, tmin

membrane (cm3/s). Gas flux was calculated from the Gfr per unit area where, gauge pressure was considered as 80-14.69 psi. Guage pressure indicates the absolute pressure difference of the feed side and permeates side (cm Hg). Then, the normalized gas flux was calculated using the eqn. 5;

7

Gas flux (5) Guage pressure and gas permeability coefficient (GPC) = NGF × membrane thickness, where, Normalized gas flux (NGF) =

permeability was expressed in barrer (1 barrer = 10-10 cm3 (STP) cm cm-2 s-1 cmHg-1).

Tensile testing: Mechanical stability of the membrane was measured using an Instron 3369 tensile tester at a strain rate of 5 mm/min at room temperature. Tensile testing was performed using a dog bone shaped sample with 20 mm gauge length, 4 mm breadth and 2.12 mm thickness prepared using Haake microinjector. Experiment was carried out twice times of each specimen for reproducibility.

Contact angle measurements: Water contact angle measurements were performed, as it provides information about the hydrophilicity of the membranes, using a Kruss F-100 tensiometer. Membrane strip of 20×10×1 mm3 size was used at room temperature in water medium. Experiment was carried out three times for every specimen for reproducibility and the average of those results is reported.

Morphological investigation: A (SUPRA 40), Zeiss SEM was used to observe the surface morphology of the membranes. A sputtering apparatus was used for gold coating on the membranes before the observation in SΕΜ.

Thermal characteristics: Thermal decomposition of PVDF and NH membranes were studied by using thermogravimetric analyzer (TGA) (Mettler-Toledo). The samples were heated from room temperature to about 600 oC at a rate of 20 oC / min under nitrogen atmosphere. The melting temperature (Tm) and heat of fusion (∆H) of the membranes were measured using DSC (Mettler 832 DSC) at a heating rate of 10 o/min. After the first melting, the samples were cooled at a constant rate of 10 o/min to understand the crystallization behavior. Further, a second heating was taken to make sure the amount of

8

crystallinity and melting temperature after removing all the thermal history in the first run. DSC was calibrated using indium.

Structural

studies:

Wide-angle

X-ray

diffraction

(Bruker

AXS

D8

Advance

diffractometer) studies of the membranes were performed with Cu Kα radiation and a graphite monochromator (wavelength, λ = 0.154 nm). The samples were scanned at 2 °/ min of diffraction angle 2θ ranging from 10 to 40o.

Electrical conductivity: The electrical dc conductivity of the membrane was measured using standard spring loaded pressure contact two-probe method across the sample. The membrane of ~130 µm thick, coated with Ag-paste for better contact, was used for the measurement. After passing constant voltage (V) across the sample from a high voltage power supply (Model: EHT-11), the current (I) was measured using a pico-ammeter (Model: DPM-111). The resistivity (ρ) was calculated from the following eqn.;

R×A (6) l where, ρ = resistivity (Ω cm), R = resistance (Ω), A = surface area (cm2) and l = thickness 1 (cm) and the conductivity ( σ ) was calculated using the relation, σ = . Each experiment ρ was carried out in triplicate and the average was taken as the conductivity of the samples. Temperature dependent conductivity was measured in a broad temperature range of 25 to 125 oC. Membrane conductivity measurement: Membrane conductivity measurements were ρ=

carried out in equilibration with water using a potentiostat/galvanostat frequency response analyzer (Auto Lab, Model PGSTAT 30). The membranes were sandwiched between two in-house made stainless steel circular electrodes (4.0 cm2). Direct current (DC) and sinusoidal alternating currents (AC) were supplied to the respective electrodes for recording the frequency at a scanning rate of 1 µA/s within a frequency range of 106 to 1 Hz. The spectrum of the blank short-circuited cell was also collected and this data was

9

subtracted (as a series circuit) from each of the recorded spectra of the membranes to eliminate cell and wiring resistances and inductances. The corrected spectra were viewed as complex impedance plots with the imaginary component of Z'' on the y-axis and the real component of Z' on the x-axis (Z = Z' - iZ''); the ionic resistance of each membrane was estimated to be the intersection of the x-axis with the extrapolation of the low frequency linear component of each plot. The membrane resistances were obtained from Nyquist plots. The proton conductivity (κm) was calculated from eqn. 7:

k m ( S / cm ) =

L ( cm )

(7)

 R ( Ω ) × A cm 2   

(

)

where, L is the distance between the electrodes used to measure the potential, R is the resistance of the membrane, and A is the surface area of the membrane. Methanol permeability measurements: Methanol permeability of the composite membranes was determined in a diaphragm diffusion cell, consisting of two compartments (80 cm3) separated by a vertical membrane with 20 cm2 effective area. The membrane was clamped between both compartments, which were stirred during the experiments. Before the experiment, membranes were equilibrated in water-methanol mixture for 12 h. Initially, one compartment (A) contained 30% methanol-water mixtures while other (B) contains double distilled water. Methanol flux arises across the membrane as a result of concentration difference between two compartments. The increase in methanol concentration with time in compartment B was monitored by measuring the refractive index using a digital refractometer (Mettler Toledo RE40D refractometer). The methanol permeability (P) finally was obtained by the equation given below: P=

1 CB(t) VBl A C A (t − t 0 ) 10

(8)

where, A is the effective membrane area, l the thickness of the membrane, CB(t) the methanol concentration in compartment B at time t, CA(t - t0) the change in the methanol concentration in compartment A between time 0 and t, and VB the volume of compartment B. All experiments were carried out at room temperature, and the uncertainty of the measured values was less than 2%. Preparation of membrane electrode assembly (MEA): MEA was fabricated by our previous technique which consists of three-layer structure (AM, anode/cathode catalyst layer and diffusion layers) [26]. The carbon paper (Toray Carbon Paper, thickness: 0.27 mm) was wet proofed with 12 wt.% PTFE solution by the brush painting method. The GDL (25 cm2 geometric area) was fabricated by coating slurry of 0.95 mg/cm2 consisting of carbon black (Vulcan XC72R) and PTFE dispersion on carbon paper. The anode was made by coating a slurry of catalyst (20 wt% Pt + 10 wt% Ru on carbon), 5 wt% Nafion ionomer solution, isopropanol, and Millipore water (catalyst ink) on GDL had a loading of 1 mg Pt and 0.5 mg Ru. While the cathode was obtained by coating the same catalyst ink lacking Ru with same loading. Electrodes were cold pressed membrane followed by curing at 60 °C for 12 h and then hot pressed at 130 °C for 3 min at 1.2 MPa. The MEA was clamped in single cell (FC25-01 DM fuel cell). The current–voltage polarization curves were recorded with the help of MTS-150 manual fuel cell test station (ElectroChemInc., USA) with controlled fuel flow, pressure and temperature regulation attached with electronic load control ECL-150 (ElectroChem Inc., USA). The measurements were performed in the air mode of operation at 10 psi pressure with 30% MeOH-water mixture at the anode side with pressure 7 psi at 70 oC for a representative membrane. 11

Results and Discussion Ionomer preparation: The functionalization and degree of sulfonation has been measured from 1H-NMR spectroscopy. Pure PVDF shows two distinguished peaks (a and b at δ = 2.23 and 2.8 ppm), attributed to head-to-head (H-H) and head-to-tail (H-T) arrangements in polymer chain, respectively [27]. Interestingly, after sulfonation reaction, the proton near the headto-head (H-H) peak (a) has split into 2.24 and 2.35 ppm peaks due to sulfonation and a/ peak intensity (δ =2.35) has increased for longer period of sulfonation time as compared to the peak at δ =2.23, suggesting greater sulfonation (Fig. 1a). Moreover, three new peaks (c, d and e) have appeared after functionalization at δ = 6.2, 7.4 and 8.1 ppm, respectively, presumably due to the attachment of sulfonate group in the PVDF chain (Fig. 1a) [28]. Nanohybrid (NH) also exhibits similar sulfonation behavior with intense peak showing greater sulfonation as compared to pure PVDF (supplementary Fig. S1). Further, peak area of the equally spaced three peaks at 7.1 ppm might be due to the aromatic proton at the chain end and their intensity increase with time owing to the greater extent of sulfonation. The percentage degree of sulfonation (DS%) has been calculated from the ratio of the integral areas of the peaks c, d, e and a, a/, b using eqn. 2 indicating considerably higher sulfonation for NH vis-à-vis PVDF (DS increase up to 27 and 32% for PVDF and NH at the reaction time of 15h, respectively) with functionalization time at a constant temperature of 65 oC (Fig. 1b) presumably due to the ease of sulfonation of γ-phase and thinner α-phase layer of crystallization in NH on top of nanoclay layers [29] against stable and thick α-phase layer in pure PVDF and will be discussed viz. structure and thermal behavior section in detail. The sulfonation reaction is occurring in both C-H and C-F 12

linkages of polymer chain and the ratio of peak integral area of c/b are 0.029, 0.032 for PVDF-s and NH-s at 15 h time while the ratio of peak integral area of e/b are 0.336 and 0.461 for PVDF-s and NH-s, respectively, for same time clearly indicate that the sulfonation has occurred most in C-F linkage as compared to C-H bond. This is to mention that c and e peaks in NMR are due to the sulfonation of C-H and C-F bond, respectively, and the intensity of both the peaks increase systematically with increasing reaction time. Based on the relative sulfonation, it is evident that C-F linkages are prone to sulfonation in comparison to C-H bonds and greater degree of sulfonation in NH as compared to pure PVDF is envisioned from the preferentially aligned polymer chains in template system in nanohybrid where fluorine atoms with δ− charge are away from the layered silicate as carbon atom having δ+ charge are nearby the negatively charged silicate layers because of electrostatic interactions (Fig. 1c). Further, higher energy state β-PVDF in nanohybrid (meta-stable state) makes it more reactive towards sulfonation than α-phase PVDF in pure polymer. FTIR studies have been performed to understand the sulfonation in the polymer chain and a new peak appears at 1036 cm-1, assigned for SO3 symmetric stretching

vibration [9] of sulfonic acid group, after sulfonation both for PVDF and NH. The intensity of the nanohybrid peak is considerably high as compared to pure PVDF suggesting better sulfonation in nanohybrid in presence of nanoclay (Fig. 1d). Further, the absorption peaks at 1500, 1640 and 1738 cm-1 after functionalization supports the sulfonation in PVDF chains [30,31]. It is worthy to mention that α-crystalline phase is noticeable for pure PVDF (764 and 795 cm-1) [32] while piezoelectric β-phase has been prominent in nanohybrid in presence of nanoclay. The peak for β-phase has been appeared 13

at 836 cm-1 both in NH and NH-s showing the presence of β-phase even after sulfonation reaction (supplementary Fig. S2) [33,34]. Figure 1e shows the UV-Vis spectra of PVDF and NH before and after the functionalization. A weak absorption band at 248 nm in NH is attributed to π→π∗ transition of olefinic bond present in layered silicate (organic modifier) nanoclay against no absorption peak for pure PVDF [35]. A strong and wide absorption band at 488 nm in PVDF-s is accredited to n→π∗ transition of sulfonate group present in polymer chain [36,37] while a significant blue shift along with greater absorption has been observed for NH-s at 450 nm mainly due to the combined π→π∗ transition of organic modifier and constricted n→π∗ transition of sulfonate group in nanohybrid [29]. It is noteworthy to mention that the above bands were not observed in unfunctionalized PVDF and NH clearly indicating sulfonation in the polymer chain. However, NMR, FTIR and UV-Vis studies confirm the sulfonation of PVDF and NH and the relative extent of sulfonation are considerably higher for NH as compared to PVDF or, in other words, functionalization can be enhanced by making composite of PVDF by dispersing suitable organically modified nanoclay in the polymer matrix followed by its functionalization.

Membrane characteristics: Water uptake is an essential criterion for any proton conducting membrane and it directly correlates the ease of proton conductivity through the membrane and mechanical strength especially toughness [38]. Figure 2a presents the water uptake of functionalized membranes which consistently increase with degree of sulfonation both for PVDF-s and NH-s reaching 27% for nanohybrid membrane as compared to PVDF-s (22%) and the result has been explained from the higher degree of sulfonation for nanohybrid vis-à-vis PVDF for a fixed duration of reaction time of 15h at 65 oC. The water uptake of Nafion is 14

38% [9] slightly higher than the present work but we restricted the degree of sulfonation up to 32% only to make a mechanically strong membrane. Ion exchange capacity (IEC) is another important parameter for a proton exchange membrane as it counts for releasing exchangeable hydrogen ions for their transport between anode to cathode. The IEC values are 0.36 and 0.50 m.mol/gm for PVDF-s and NH-s, respectively, indicating the presence of large number of exchangeable proton in nanohybrid membrane as compared to PVDF membrane primarily due to greater number of attached sulfonate groups present in nanohybrid membrane. Moreover, it is expected that chemically tagged sulfonated group would alter the surface properties including wettability of the membranes. Contact angle measurement of the developed membranes shows systematic decrease with increasing degree of sufonation or reaction time both for PVDF and NH (Fig. 2b). The contact angle is considerably less for NH-s as compared to PVDF-s suggesting better wettability for nanohybrid membrane against PVDF membrane. Initial higher contact angle (78±2o) of nanohybrid as compared to pure PVDF (75±3o) is due to slight hydrophobic nature of nanohybrid in presence of organically modified nanoclay while the greater degree of sulfonation in nanohybrid make the membrane more hydrophilic in nature resulting lowering of contact angle as compared to PVDF. However, the decrease of 26 degree (78→52o) of contact angle has been reported for NH-s membrane against 18 degree decrease (75→57o) observed in PVDF clearly indicate better hydrophilicity in nanohybrid membrane as compared to pure polymer for a fixed sulfonation time of 15h. Similar hydrophilic nature of membranes due to functionalization of the polymer matrix is reported in the literature [39,40]. To understand the molecular size of the polymer chain due to functionalization, gel permeation chromatography (GPC) has been performed to discern the hydrodynamic 15

volume of the sulfonated species and their comparison. GPC traces of pure PVDF and sulfonated species for various times are presented in Fig. 2c, showing significant shift of elution time with increasing reaction time. The lower elution time for greater sulfonated species is a clear indication of higher molecular weight or larger hydrodynamic volume resulting from direct attachment of sulfonate group in the polymer chain. The weight average molecular weight of pure PVDF is 1.04×105 (Mn) while using the same calibration curve, the molecular weight is around 22.8×105 for PVDF-s at 15 h (taking the peak value of the chromatogram) of the sulfonated species, which is unexpectedly on the higher side. Hence, unusually higher hydrodynamic volume of sulfonated species is presumably due to ionic cluster/association formation in the solution phase and the fact of further lowering of elution time for sulfonated nanohybrid auxiliary strengthen the clustering phenomena arising from greater amount of sulfonate group (higher degree of sulfonation) in nanohybrid (Fig. 2d). However, bimodal distribution is evident in the chromatogram of functionalized specimens and the peak position corresponding to higher elution time is probably the molecular weight of sulfonated species. From the above peak, the molecular weight of functionalized nanohybrid is clearly manifested as compared to PVDF-s from the lower elution time of NH-s (Fig. 2c &d). Another essential property of the membrane is its permeability and the polymeric membrane should not permeate the gas easily. Comparative gas permeation has been shown in Fig. 3a indicating considerably lower permeation after the functionalization both for PVDF and NH, while amongst functionalized specimens, NH-s shows the highest barrier (63 barrer) property mainly because of strong ionic cluster present in sulfonated nanohybrid as compared to sulfonated PVDF (82 barrer) as discussed above. The higher barrier property of NH (105 barrer) vis-à-vis PVDF (140 barrer) is primarily due to the 16

dispersed two dimensional nanoclay in the polymer matrix, by creating a tortuous path. Incorporation of nanoparticle into the polymer matrix usually reduces the permeation rate of gas [41-43] while ionic clusters leading to compact structured material is the unique phenomena of functionalized polymeric membrane especially in presence of nanoparticle. It is also expected that direct fuel permeability would be restricted in this functionalized nanohybrid membrane leading to better fuel efficient membrane as fuel permeability and proton transport are the main criteria of any membrane [44,45]. An ideal fuel cell membrane should be a free standing film with sufficiently mechanically stable to withstand the high osmotic pressure [46]. The stiffness and toughness of the functionalized membranes has been verified using universal testing machine under a constant stretching rate. The stress-strain curves of the membranes are presented in Fig. 3b showing much higher elongation at break for nanohybrid membrane (66%) as compared to PVDF membrane (22%). The toughness values, measured from the area under stress-strain curve, of the membranes are 1.18 and 2.19 kJ.m-3 for PVDF-s and NH-s, respectively, indicating superior toughness for nanohybrid membrane (86% higher) presumably due to structural change of the matrix (facile form β-phase) in presence of nanoclay. Elastic moduli of PVDF-s and NH-s membranes are 902 and 715 MPa, respectively, whereas the modulus of neat PVDF is 720 MPa [29]. Surface morphologies of the membranes, before and after functionalization, are presented in Fig. 3c. Pure PVDF has shown nice spherulitic pattern while nanohybrid exhibit predominantly mesh-like structure because of structural change (α→β phase conversion) [35,36]. White spots (roughening) are observed both in PVDF and its nanohybrid membrane after sulfonation due to ionomer formation, eliminating the special feature of pure PVDF and NH. Larger number of white spots in nanohybrid membrane is 17

noticed as opposed to PVDF membrane again indicate greater functionalization in nanohybrid. Good dispersion of nanoclay is evident in NH in the form of white spots while clear micrograph is noticed in pure PVDF.

Thermal behavior The

thermal

stability

of

the

membranes

was

investigated

by

using

thermogravimetric analyzer (TGA) and the weight loss as a function of temperature has been shown in Fig. 4a. Pure PVDF degrades at around 475 oC (5 % weight loss) [47] while with increasing reaction time / degree of sulfonation, the degradation temperature gradually decrease to 266 and 240 oC for 12 and 15 h reaction time, respectively, showing lowering of degradation temperature of the sulfonated species. Functionalized membranes exhibit two stages of degradation and the initial degradation is due to the dissociation of sulfonate group (–SO3H) [26] followed by the degradation of main chain at higher temperature against the sharp one stage degradation before functionalization. Functionalized nanohybrids exhibit similar behavior while the proportion of the first phase of degradation is relatively high as compared to PVDF presumably due to greater amount of sulfonation in nanohybrid (supplementary Fig. S3). Pure nanohybrid shows the first degradation point at 390 oC whereas, the degradation of sulfonate group takes place at 235 and 169 oC for functionalized nanohybrids with the reaction time of 12 and 15h, respectively. The melting behavior of different sulfonated PVDF has been presented in Fig. 4b showing decreasing melting temperature (Tm) with increasing degree of sulfonation / time of PVDF and this decrease of Tm by ~11 oC is caused by the insertion of sulfonate group in the main chain which inhibit the crystallization in presence of bulky –SO3H group 18

[48,49]. Pure nanohybrid shows slight higher melting temperature of 174.5 oC as compared to neat PVDF indicating the presence of γ-phase along with β-phase induced by the nanoclay platelets (Fig. 4c) [35]. On contrary to PVDF, the melting temperature considerably increases for nanohybrids by ~2 oC after sulfonation which systematically varies on reaction time presumably due to the development of metastable phase which appears because of distorted α- or γ-phase upon crystallization on the surface of the nanoclay platelets [29] and will be discussed later in details. This metastable phase basically facilitate the higher ionomer formation in nanohybrid with respect to PVDF. Further, PVDF exhibit spherulitic structure with double melting endothermic peaks but, gradually the first melting peak disappears with higher sulfonation while mesh-like crystallites are evident in nanohybrid (Fig. 3c) showing only one melting peak [50]. Moreover, the heat of fusion (∆H) of neat PVDF and NH are 61.5 and 39.2 J.g-1, respectively, while it decrease significantly with increasing degree of sulfonation or reaction time both for PVDF-s and NH-s. The lesser value of ∆H in unfunctionalized nanohybrid is due to the formation of β-phase in presence of nanoclay while the significant decrease in functionalized nanohybrid as compared to PVDF is mainly due to greater degree of sulfonation in naohybrid vis-à-vis PVDF (Fig. 4d). The heat of fusion (∆H) decreases drastically from 61.5 to 44 J.g-1 and 39.2 to 15.3 J.g-1 after sulfonation for pure PVDF-s and NH-s membrane, respectively, by reducing the chances of fitting the substituted –SO3H bulky group into crystalline lattice. Structural alteration The nanostructural investigation of the nanohybrids, before and after sulfonation, has been presented in Fig. 5a. The XRD peak of nanohybrid shows at 2θ = 2.8o, subsequent to d001 basal spacing of 3.16 nm, in PVDF is attributed to intercalation of polymer inside the nanoclay galleries [23,24,51]. It is to mention that pure PVDF does not show any nanostructure. The intercalated structure becomes exfoliated and/or disordered after 19

functionalization primarily due to the additional bulky sulfonate group insertion into the polymer chain which pushed out the gallery further. The peak at 2θ = 6.1o is appeared from sandwiched structure arising from the crystallization of polymer on both sides of the silicate layers [22,29]. Figure 5b shows the crystalline structure of PVDF and PVDF-s membranes functionalized for various time. The peaks at 17.6o, 18.3o, 19.9o and 26.6o are corresponding to the α-phase [35] and the structural integrity is maintained even after sulfonation for longer period of time. In contrast, PVDF crystallizes in piezoelectric βphase in nanohybrid in presence of nanoclay as evident from the peak position at 2θ = 20.65o (corresponding to (200/110) planes) and another peak at 2θ = 19o which appears due to metastable phase (β- or γ-conformation) in Fig. 6c [34]. The intensity of β-peak consistently increases with degree of sulfonation of nanohybrid as clear from the deconvoluted peaks as shown in Fig. 5d. Fraction of β-content was calculated from area under respective peak which show 24, 28, 36 and 42% of β-phase for NH, NH-s of 3, 9, and 15h reaction time, respectively. Further, α−, β− and γ−phase contents of pure NH are 40, 24 and 4%, respectively, while the similar phases appear at 23, 42 and 8% in NH-s (15 h sulfonation time). Now, it is apparent that there is no structural alternation for PVDF due to sulfonation while significant improvement in β-phase content occurs in nanohybrid for similar reaction condition. During sulfonation of solid film in reaction condition, it is difficult to visualize that crystallization will occur during reaction or reorganization of crystalline phase at so low temperature (65 oC). Rather, the sulfonation on α-phase covering the islanded β-phase crystalline structure on the surfaces of nanoclay is envisioned which amorphizes some part of α-phase and as a result overall β-phase increases in the system [52]. A model has been proposed in Scheme I, where β-phase (all trans) is crystallized on top of nanoclay [53] which gradually change into α-phase ( ¹

TGT G ) as the crystalline phase is away from nanoclay surface through an intermediate γ¹

phase (T3GT3 G ) , and thereby form an island structure whose surface is wrapped up with α-phase. Those α-phase PVDF chain is susceptible to sulfonation causing considerable amorphization as evident from the significant lowering for heat of fusion for NH-s as compared to PVDF-s (Fig. 4d). Electrical conductivity: So far, we have demonstrated the functionalization of PVDF and its nanohybrid and their relative improvement in various properties showing more sulfonation in nanohybrid as compared to pure PVDF. In order to be used as fuel cell membrane, the film should be conducting and the bulk electrical dc conductivity of PVDF and nanohybrid membrane in the film form was measured. The conductivity of pure PVDF membrane after sulfonation at 65 oC for 3h shows a jump from 10-15 to 4.8×10-9 S.cm-1 and subsequent increase of sulfonation time improves the conductivity further to 5.8×10-4 S.cm-1 (Fig. 6a). Similar behavior was observed for nanohybrid but the relative improvement in dc conductivity is considerably higher in nanohybrid (5×10-3 S.cm-1) vis-à-vis PVDF in similar condition mainly due to greater extent of sulfonation in nanohybrid as reported above. Further, increase of sulfonation temperature for a fixed time of 4h also exhibit increase of dc conductivity both for PVDF-s and NH-s while the relative increment is higher in NH-s as compared to PVDF-s (Fig. 6b). Hence, the nanohybrid membranes demonstrate higher conductivity by more than one or two orders of magnitude as compared to PVDF reaching 20

the semi-conductor level starting from an insulating behavior of nonfunctionalized PVDF/NH. The conductivity was also measured at different temperatures to calculate the E activation energy (Ea) following the Arrhenius equation, ln σ = ln σ 0 − a where, σ is RT -1 -1 the conductivity of the samples (S.cm ), σo, pre-exponential factor (S.K .cm-1), Ea, activation energy of conduction (kJ.mol-1) governed by the Grotthuss type mechanism [54], R, ideal gas constant and T, temperature (K) (Fig. 7c). The neat PVDF shows the activation energy of 10.6 kJ.mol-1 while the value reduces in nanohybrid to 6.8 kJ.mol-1. However, the activation energy remains same for PVDF after sulfonation while significant enhancement has been observed for NH-s indicating the stability of membrane conduction at higher temperature and the value is equivalent to literature reported value [55]. The activation energies (Ea) of nanohybrid and PVDF membranes are shown in Table 1.

Proton conductivity and methanol permeability: The proton conductivity and methanol permeability are the two major important parameters for any fuel cell membrane. Higher proton conductivity and lower methanol permeability values are preferable for a good membrane. Proton conductivity and methanol permeability of functionalized PVDF and nanohybrid membranes are presented in Table 2 and compared the values of standard Nafion 117 membrane. Functionalized PVDF and NH membranes show high proton conductivity as compared to pure PVDF and non functionalized NH and, amongst the functionalized membranes, NH-s shows highest conductivity (6.13×10-2 S.cm-2), the value close to Nafion 117 (9.56×10-2 S.cm-2) presumably due to the greater degree of sulfonation in nanohybrid as compared to PVDF as described earlier. Proton conductivity is closely related to water uptake and ion exchange capacity of any membrane and have more favorable condition for nanohybrid vis-à-vis PVDF and thereby, explain the higher conductivity of NH-s. Ladewig et al. showed that conductivity of Nafion nanohybrid membrane is lower than pure membrane due to the presence of inorganic agglomerates and particles which alter the proton-transfer pathway [45]. In contrast, two dimensional nanoplatelets alter the structure and thereby 21

degree of sulfonation and help improving proton conductivity in this work. Proton conductivities at higher temperature are shown in Fig. 7a. Interestingly, the slope value is considerably becoming lower in nanohybrid which further reduced in NH-s but display higher values of absolute conductivities. The values of the slope indicate the activation energy (Ea), minimum energy required for proton transport phenomena across the membrane, which decrease in nanohybrid (5.75 kJ mol-1) as compared to pure PVDF (16.59 kJ mol-1) and further reduced in NH-s (4.92 kJ mol-1). This value is even lower than Nafion 117 (6.52 kJ mol-1). Usually, proton transport occurs across the membrane by two mechanisms, i) a vehicle mechanism, where proton combines with water molecule forming H3O+ (hydronium ion) which then diffuse through membrane and, ii) Grotthus mechanism, when proton hopps through the ionomer chains having sulfonate group (-SO31

)

attached to PVDF main chain (Fig. 1a). However, the higher conductivity in

functionalized nanohybrid (NH-s) is envisaged in the course of frequent availability of sulfonate group in PVDF chain through which either hopping or diffusion processes of proton are made easier as compared to PVDF-s where the gap between two sulfonate groups in a polymer chain is considerably large (lower degree of sulfonation). Moreover, the activation energy of proton conduction for pure PVDF is within normal limit of 14 40 kJ mol-1 while the complex conduction mechanism in nanohybrid and functionalized membrane makes significantly lower activation energy [56]. It is needless to mention that the activation energy also decreases with increasing ion exchange capacity (IEC) of the membrane [9]. Methanol permeabilities through different membranes are presented in Table 2 showing relatively lower permeabilities. However, methanol cross over rate of PVDF based membrane (PVDF-s and NH-s) is two order lesser than Nafion 117 membrane. The 22

significantly less methanol cross over (less permeability) of functionalized PVDF membrane as compared to standard Nafion lies in the superior crystalline phase along with compact spherulitic morphology of PVDF. The aggregate pores, responsible for mass transport (methanol), becomes considerably less around polymer chains in PVDF based membrane. Significant low methanol permeation of functionalized membrane is visualized further from strong ionic cluster of sulfonated species (both for PVDF-s and NH-s) as discussed in GPC and gas permeation section above. However, the importance of nanohybrid based membrane against Nafion is revealed through high proton conduction and significantly low methanol permeability. For comparison of the applicability of DMFCs, proton conductivity and methanol permeability were used to calculate the selectivity parameter (SP) following the Equn.: Km , where, PMeOH is the methanol permeability (cm2 s-1). The selectivity data has PMeOH been presented in bar diagram (Fig. 7b) showing much higher value of NH-s as compared to PVDF-s. Further, the SP value enhances in nanohybrid presumably due to higher conductivity and lower permeability. Interestingly, the SP values of functionalized PVDF membranes are very high (even in comparison to Nafion 117) indicating the much better performance of the developed membranes. These results can be attributed to the lack of interaction between methanol and ionomers (sulfonate group attached PVDF chain). It is also noticed that SP values of the functionalized membranes increased with increasing the operating temperature, which was significantly higher than that of Nafion 117 (0.8×105 S.s cm-3) [56]. However, higher SP values at high temperature of the developed membranes indicate great advantage of functionalized nanohybrid membrane (NH-s) especially for high temperature applications. Fuel cell performance in DMFC: SP =

The performance of the functionalized membranes was measured by recording I-V polarization curves in DMFC. Figure 8a has shown the fuel cell performance (polarization curve) of PVDF-s and NH-s membranes indicating slow decay of cell voltage with increase in current density. NH-s shows the best performance as evident from the lower slope of I-V curve. PVDF-s and NH-s show higher potential throughout the current density 23

measured here. PVDF-s shows lower slope up to 50 mA / cm2 followed by a sharp decrease while NH-s exhibit lower slope up to 80 mA / cm2. Further, maximum power density is obtained in case of NH-s (Fig. 8b) suggesting functionalized nanohybrid membrane is superior to the others. It is noteworthy to mention that both cell voltage including lower slope of I-V curve and power density of NH-s is considerably enhances vis-à-vis standard Nafion 117 membrane and, thereby, demonstrate an improved fuel cell membrane. The high cell voltage and power density of functionalized membranes are envisioned from the high proton conduction and low methanol permeability of the ionomer. Every PEM membrane suitable for fuel cell application should increase in current and power density which in turn depends on the methanol concentration fed to the anode [56]. Nafion 117 membrane has shown the power density of 9.28 mW / cm2 at a current density of 45.1 mA / cm2 while NH-s in this study exhibits a maximum of 33 mW / cm2 current density [26]. The excellent fuel cell performance of functionalized PVDF nanohybrid membrane as compared to Nafion117 is presumably due to enhanced sulfonation in presence of two dimensional layered silicate present in the nanohybrid which exhibit significantly lower methanol permeation along with mechanically tougher by design. New functionalized membrane shows the new generation smart fuel cell membrane using common thermoplastics like PVDF.

Conclusions: Poly(vinylidene fluoride) nanohybrid has been prepared with organically modified two-dimensional layered silicate followed by its sulfonation using chlorosulfonic acid at control condition to fabricate fuel cell membrane. The degree of sulfonation has been varied by changing the sulfonation time while the extent and location of functionalization 24

in polymer chain has been verified through NMR, FTIR and UV-Vis studies. The greater sulfonation in nanohybrid as compared to pure PVDF is explained from the specific alignment of polymer chain on top of layered silicate template. Required criteria like water uptake, ion exchange capacity, contact angle (measure of hydrophilicity) and permeability for fuel cell membrane exhibit favorable values and indicate superior assessment for functionalized nanohybrid as compared to sulfonated PVDF. PVDF crystallizes in piezoelectric β−phase in presence of nanoplatelets and the extent of β−phase has increased further after functionalization of nanohybrid which has been verified again through FTIR and DSC measurement leading to propose a model for crystallization behavior. Toughness of functional nanohybrid membrane has shown to be improved as compared to sulfonated PVDF and, thereby, provide a membrane with tougher by design induced by nanoplatelets. DC conductivity of the functionalized membrane in solid state has shown to be in the semiconducting range against completely insulating behavior of pure PVDF. Proton conductivity of the sulfonated nanohybrid is comparable to the value of commonly used Nafion membrane with considerably lower activation energy. Methanol crossover is greatly reduced in sulfonated nanohybrid membrane by two orders of magnitude as compared to Nafion. Membrane electrode assembly (MEA) has been constructed using various membranes and the potential is shown to be considerably higher at lower current density with lesser slope value indicating better performance of the functionalized nanohybrid membrane vis-à-vis sulfonated PVDF. Power density of sulfonated nanohybrid membrane exhibits significantly higher value of 33 mW/cm2, against the value of 11 mW/cm2 of standard Nafion at similar current density, certainly demonstrates the superior membrane with added possibility of smart membrane arising from the piezoelectric nature of the membrane. 25

Acknowledgements The author (Karun Kumar Jana) thank for award of Senior Research Fellowship of CSIR, India.We gratefully acknowledge the help from Dr. D. K. Avasthi, Dr. Pawan K. Kulriya, and Prof. Nira Misra for support in XRD and contact angle measurements.

Supporting Information Available: Figures S1-S5 and Table S1.

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Table 1: Activation energy of PVDF and NH membranes for different sulfonation time

Membrane Activation Energy (Ea) in kJ.mol-1 for different sulfonation time

PVDF NH

0h 10.65 6.88

3h 8.78 7.97

9h 10.72 13.77

12h

15h

11.21 14.35

12.49

18.14

Table 2: Membrane conductivity (κm), methanol permeability (P) (at 30 oC) and energy of activation (Ea) values for different membranes measured at 30 OC with 30% MeOH-water mixture. The values of Nafion 117 are taken from Ref. 1.

Membrane PVDF-s NH-s Nafion117*

κm (10-2 S cm-1) 3.62 6.13 9.56

P (10-8 cm2 s-1) 3.14 3.21 131.0

34

Ea (kJ mol-1) 8.87 4.92 6.52

Amorphous TGTG T3GT3G TTTT Clay Functionalization

Scheme I: Schematic representation of the structural change over showing less αcrystalline phase after functionalization of nanohybrid which gets converted into amorphous phase.

35

b

a/ a

e d

(a)

c

15h

b Hb SO3He (C

CF

Ha/ Ha H

CH2

SO3Hc

(b)

e

CF2 CF2 C

d

C a )n CH

a/ a

SO3Hd

c

9h

b(H-T)

Hb

Ha

( CHb CF2

CF2

CH2

CF2

a(H-H)

CH2 CHa )n

PVDF

(c)

(d)

H

+

δ δ−

Reflectance / a.u.

δ−

F

H H

F

δ−

F

δ+ δ−

1650 NH-s 1734

H H

1500

1036

NH

PVDF-s PVDF

F

H 1800

1500

1200

900 -1

Wavenumber / cm

(e)

0.6

450

NH-s NH PVDF-s PVDF

248

Reflectance / a.u.

488

0.4

0.2

0.0 200

400

600

800

Wavelength / nm

36

Figure 1: (a) 300 MHz 1H NMR spectra of neat PVDF and PVDF-s for 9 and 15h sulfonation. The positions of the proton NMR peaks have been shown in the inset, (b) Degree of sulfonation plot for PVDF and NH as a function of sulfonation time (0, 3, 9, 12 and 15h). (c) Schematic representation of chain alignment on layered silicate showing facilitation of sulfonation in C C-F bond as compared to C-H H linkage in nanohybrid, (d) FTIR spectra of pure PVDF PVDF, PVDF-s, NH and NH-s. The vertical lines represent the position of extra peaks due to sulfonation, (e) UV-Vis absorption spectra of pure PVDF, PVDF-s, NH and NH-s (sulfonation time for 15 h at 65 oC).

(a)

(b)

(c)

(d) 15h

15h 12h

12h 9h

9h

NH

PVDF

18

21

24

27

18

21

24

t / min

t / min

37

27

Figure 2: (a) Water ater uptake of PVDF-s and NH-s membranes of different sulfonation time (0-15 h), (b) water contact angles for the membranes as a function of sulfonation time, (c) GPC traces for pure PVDF and its sulfonated species for different time as mentioned, (d) GPC traces for pure nanohybrid hybrid ((NH) and its sulfonated species for different time as mentioned.

(a)

(b)

σ / MPa

45

30 PVDF-s

NH-s

15

0

(c)

PVDF

PVDF-s

NH

NH-s

0

15

30

ε/%

45

60

Figure 3: (a) Permeability of N2 gas through PVDF and NH membranes (before and after functionalization). (b) stress--strain curves for PVDF and its nanohybrid membranes. (c) SEM images of pure PVDF,, pure NH, PVDF-s and NH-s (15 h sulfonation).

38

(a)

(b)

1.0

15h

Heat flow / a.u.

Weight fraction

12h 0.8

0.6 15h 12h 3h PVDF

0.4

163.3

9h 169.4

3h

172.7

PVDF

173.8

174

0.2

100

200

300

400

140

500

Temp / C

(c)

160

170

180

190

(d)

15h

Heat flow / a.u.

150

T / °C

0

12h

176.1 175.4

9h 3h

175.2

NH

175

174.5

140

150

160

170

180

190

T / °C

Figure 4: (a) TGA thermograms of neat PVDF and PVDF-s for different sulfonation time, (b) DSC thermograms of pure PVDF and time dependent sulfonated PVDF (c) DSC thermograms of nanohybrid and its sulfonated species (NH-s)) for different time. The vertical lines represent the melting point of the corresponding specimens. (d) Heat of fusion as a time of sulfonation time for pure PVDF, NH and their sulfonated species. 39

(b)

Clay

2

4

6

8

15h

Intensity / a.u.

Intensity / a.u.

Intensity / a.u.

(a)

10

2θ / deg

15h 9h

9h

3h

3h PVDF

NH

2

4

6

8

10

12

2θ / deg

18

24

30

36

2θ / deg

(d)

15h

Intensity / a.u.

Intensity / a.u.

(c) 15h 9h

3h

NH

NH

16

18

20

22

16

24

18

20

22

24

2θ / deg

2θ / deg

Figure 5: (a) XRD patterns of layered silicate clay and different NH-s as indicated. (b) XRD patterns of pure PVDF and PVDF-s for different time as indicated showing α-phase crystalline state, (c) XRD patterns of NH and NH-s for different time as indicated showing

β-phase crystalline state, the vertical lines indicate β-phase peak position, 40

(d)

Deconvoluted XRD patterns of NH and representative NH-s showing larger peak area corresponding to β-phase.

(a)

-4

(b)

o

T = 65 C

10

t = 4h -6

-4

10

10

-8

10

-1

σ / s.cm

-1

σ / s.cm

-8

10

-10

10

-12

10

-12

10

NH-s PVDF-s

10

15

10

NH-s PVDF-s

-14

-16

10

0

3

9

12

-16

t/h

(c)

0

40

60

80

100

T / 0C NH-15h NH-3h PVDF-15h PVDF-3h

-3

ln σ

-6

-9

-12

-15 2.4

2.6

2.8

3.0

3.2

3.4

-1

1000/T (K )

Figure 6: Electrical dc conductivity measurements of (a) Time dependent measurement of PVDF-s and NH-s. The horizontal lines show the respective blank (nonfunctionalized PVDF and NH) (b) Temperature dependent conductivity of PVDF-s and NH-s. The horizontal lines show the respective blank (nonfunctionalized PVDF and NH), (c)

41

Comparison of electrical conductivity as a function of temperature (1/T) ooff various degree of sulphonation / time of PVDF and NH membranes as mentioned.

(a)

(b)

PVDF-s NH-s

16 12

5

SP X 10 / S s cm

-3

20

8

4 0

Figure 7: a) Arrhenius plots of ln km vs. 1/T for PVDF-s and NH-s, b) SP values of different membranes as mentioned measured at room temperature temperature. SP value of standard Nafion 117 is 0.75×105 S.s cm-3 at room temperature.

42

(a)

(b)

Figure 8: Direct methanol fuel cell performance (current (current-voltage voltage polarization studies) of PVDF-s and NH-s (15h sulfonation).

43

Table of Content Functionalized Poly(vinylidene fluoride) Nanohybrid for Superior Fuel Cell Membrane

Karun Kumar Jana1, Chumki Charan2, Vinod K. Shahi2, Kheyanath Mitra3, Biswajit

H δ−

H

F δ+

H

δ−

F− δ−

H

F δ+ δ

H H

δ−

F

Power density / (mW / cm2)

Ray3, Dipak Rana4, and Pralay Maiti

*1

H

30

δ−

Nanohybrid H

20

F δ+

H

Nafion

10

SO3H

δ−

H

F δ+ δ

0 0 20 40 60 802 Current density / (mA / cm )

Functionalization

clay TTTT T3GT3G TGTG

44

H H

SO3H

Highlights •

Light weight polymeric hybrid and piezoelectric membrane for fuel cell.

•

MEA show high proton conductivity, high potential at low current density with significantly high power density.

•

Extremely low methanol crossover vis-à-vis Nafion along with toughened nature, effective for next generation fuel cell.

•

The cause of higher efficiency in nanohybrid has been worked out.

•

The structural development in presence of nanoparticle has been looked into shading light on the functionalization.

45